Method of depositing oxide passivation layers on high temperature superconductors

ABSTRACT

The present invention provides a method of depositing a passivation layer on the surface of a superconducting ceramic oxide wherein the passivation layer is a layer of an oxide of Al, Bi, Si or Al-W.

GRANT INFORMATION

The present invention was made with the support of the Office of Naval Research under contract No. N0014-86-K-0427. The Government has certain rights in this invention

This is a continuation of application Ser. No. 07/554,016, filed Jul. 17, 1990, which was abandoned upon the filing hereof, which is in turn a division of Ser. No. 7/246,100 filed Sep. 9, 1988, now U.S. Pat. No. 4,965,244.

BACKGROUND OF THE INVENTION

Superconductors are materials which can conduct electricity with virtually no resistance when they are maintained at a certain temperature, referred to as the superconductive transition temperature (TC). For example, pure metals or alloys such as niobium-tin (Nb₃ Sn) reach the superconductive state when they are cooled below 23 K. That degree of cooling requires the use of liquid helium, which condenses at 4 K. Liquid helium is expensive and is difficult to manipulate.

A major breakthrough in the commercial development of this technology cam in January of 1987, when a yttrium-barium copper oxide ceramic, reported to be Y₁.2 Ba₀.8 CuO₄, was prepared which achieved superconductivity at a Tc of about 90 K. See M. K. Wu et al., Phys. Rev. Lett., 58, 908 (1987). This degree of cooling can be readily acomplished with liquid nitrogen (boiling point, bp, 77 K. or -196° C.), which is much less expensive and easier to handle than is liquid helium (bp 4.2 K.).

The advent of liquid nitrogen-cooled superconductors could be a boon to utilities, industry, electronics, transportation and medicine. For example, power companies envision superconductive transmission lines, buried underground, that would carry current with no dissipative losses or generation of heat. Superconducting devices could lead to smaller, more powerful supercomputers. Because these chips produce no waste heat, they could be packed closer together, allowing the size of electronic boxes to be reduced. This size reduction means that signals would take less time to travel between switching devices and circuit elements, leading to smaller, faster computers.

Since the discovery of high temperature superconductivity in the 2-1-4 oxides (of the form La_(2-x) ^(Sr) _(x) CuO₄), the 1-2-3 oxides (of the form YBa₂ Cu₃ O_(7-x)), the 2-1-2-2 oxides of Bi and Tl (of the form Bi₂ Ca_(1+x) Sr_(2-x) Cu₂ O_(8+y) or Tl₂ Ca_(1+x) Cu₂ O_(8+y)), and the 2-2-2-3 oxides of Bi and Tl (of the form Bi₂ Ca_(2+x) Sr_(2-x) Cu₃ O_(10+y) or Tl₂ Ca_(2+x) Ba_(2-x) Cu₃ O_(10+y)) it has been apparent that the incorporation of these materials into existing and new technologies will require the solution of a large number of materials-related problems. Analogous problems can be anticipated for all such CuO-based superconductors, including systems presently in the early stages of development.

First, there are issues related to materials synthesis so that structures can be fabricated with predetermined shapes, sizes and current-carrying ability. These range from macroscopic to microscopic. Second, there are challenges related to the fabrication of superconducting thin films on a variety of substrates, with Si being an obvious choice from the perspective of microelectronic devices. Third, there are issues related to the formation of stable ohmic contacts, particularly for small samples and thin films. Fourth, there are problems related to the passivation, protection or encapsulation of small structures such as fibers or thin films, so that the superconducting oxides can be used under a wide range of environments.

Early work has shown that the surfaces of these high-Tc ceramics are highly reactive. In particular, adatoms which are reactive with respect to oxide formation cause disruption of the 2-1-4, 1-2-3, and 2-1-2-2 surfaces by withdrawing oxygen from the lattice to form undesirable oxide overlayers with thicknesses that are probably kinetically limited. In a broad sense, passivation layers act to prevent molecular motion across the layer, providing a stabilizing environment which protects the surface of the superconductor from the ambient atmosphere. More specifically, the passivation layers can serve as barriers against loss of oxygen from the superconductor, can isolate one device from another, can act as components in the semiconductor structure, or can provide electrical isolation of multilevel conductive systems.

H. M. Meyer et al., in Appl. Phys. Lett., 51, 1118 (October 1987) reported that evaporated gold films can passivate La₁.85 Sr₀.15 CuO₄ superconductor surfaces against chemical attack. In an attempt to coat super-thin films with dielectric films, Y. Ichikawa et al., in J. Appl. Phys., 27, L381 (1988) used rf-magnetron sputtering to deposit Nb₂ O₅ films and Al₂ O₃ films on films of the Y-Ba-Cu-O superconductive ceramic disclosed by M. K. Wu et al., cited hereinabove. The thickness of the dielectric films was reported to be from 600-1350Å. However, these layers deleteriously modified the crystal structure of the Y-Ba-Cu-O films, leading to a broadened superconducting transition zone and a lowered Tc. This group also reported Ba atom diffusion into the dielectric layer, indicating substantial disruption of the ceramic lattice.

Therefore, a need exists for methods to apply stable, effective passivation layers to superconductive ceramic oxides which do not disrupt the useful properties of the superconductor.

SUMMARY OF THE INVENTION

The present invention provides a method to apply a passivation layer to the surface of a superconductive ceramic oxide. The method comprises evaporating aluminum (Al), bismuth (Bi), silicon (Si) or both aluminum and tungsten (Al-W) in an atmosphere of thermally or chemically activated oxygen (O*) so that a dielectric layer of an oxide of Al, Bi, Si or Al-W is applied to said surface. The present invention also comprises applying a passivation layer of CaF₂ to the surface of a superconductive ceramic oxide by molecular beam deposition of CaF₂ thereon, which deposition does not require the use of activated oxygen. The passivation layers applied are dielectric, as opposed to conductive pure metal layers employed to form ohmic contacts. The present oxide or fluoride passivating layers do not substantially disrupt the ceramic lattice of the superconductor, and therefore can limit the degradation of its electric or physical properties.

Although the present method is exemplified primarily by the deposition of thin (about 12-100Å) layers to facilitate study of interface-overlayer interactions, layers of any appropriate thickness can be applied. Perferably, the present passivation layers are at least about 50Å in thickness, most preferably about 100Å to 5,000-10,000Å in thickness.

The present method can be employed to apply a passivation layer to the surface of any of the known superconductive ceramic oxides. The structures of several of some preferred classes of these materials is summarized in Table I, below. Other copperoxide-based superconductive ceramic oxides which can be used as substrates in the present method are described hereinbelow.

                  TABLE I                                                          ______________________________________                                         Superconductive Ceramic Oxides                                                 Formula.sup.6     Abbreviation                                                 ______________________________________                                         La.sub.2-x A.sub.x CuO.sub.4                                                                     2-1-4.sup.1,.sup.4                                           RBa.sub.2 Cu.sub.3 O.sub.7-x                                                                     1-2-3.sup.2,.sup.5                                           Bi.sub.2 Ca.sub.1+x Sr.sub.2-x Cu.sub.2 O.sub.8+y                                                2-1-2-2.sup.3                                                Bi.sub.2 Ca.sub.2+x Sr.sub.2-x Cu.sub.3 O.sub.10+y                                               2-2-2-3                                                      Tl.sub.2 Ca.sub.1+x Ba.sub.2-x Cu.sub.2 O.sub.8+y                                                2-1-2-2                                                      Tl.sub.2 Ca.sub.2+x Ba.sub.2-x Cu.sub.3 O.sub.10+y                                               2-2-2-3                                                      ______________________________________                                          .sup.1 J. G. Bednorz et al., Z. Phys. B., 64, 189 (1986), (A = Sr).            .sup.2 C. W. Chu et al., Phys. Rev. Lett., 58, 405 (1987), (R = Y).            .sup.3 H. Maeda et al., Jpn. J. Appl. Phys., 27, L209 (1988); Z. Z. Sheng      et al., Nature, 332, 55 (1988).                                                .sup.4 A = Ba, Sr, Ca.                                                         .sup.5 R = lanthanide element, e.g., Y, Sm, Eu, Gd, Dy, Ho, Yb.                .sup.6 X and y are stoichiometric values required to satisfy the chemical      valence states necessary for superconductivity.                          

The present invention is also directed to a shaped body of a superconductive ceramic oxide, such as a wire, a microelectronic device, film or the like, having at least one surface thereof coated with a passivation layer of (a) an oxide of Al, Bi, Si or Al-W or (b) CaF₂.

As used herein with respect to the present oxide fluoride layers, the term "passivation" is intended to encompass the protection and/or encapsulation of superconductive structures such as fibers or thin films, as well as their electrical isolation. Therefore, the present invention also comprises a shaped body of a superconductive ceramic oxide having at least one surface coated with a passivation layer of (a) an oxide of Al, Bi, Si or Al-W, or (b) CaF₂.

BRIEF DESCRIPTION OF THE FIGURES

FIG. 1 is a graphical depiction of the Cu 2p_(3/2) core level emission for a 1-2-3 superconductor surface treated in accord with the present method.

FIG. 2 is a graphical depiction of the Cu 2p_(3/2) core level emission for a 1-2-3 superconductor surface treated in accord with the present method.

FIG. 3 is a graphical depiction of the O is core level emission for a 1-2-3 superconductor treated in accord with the present method.

FIG. 4 is a graphical depiction of the Cu 2p_(3/2) core level emission of a 1-2-3 superconductor treated with Al or Al in thermally or chemically activated oxygen (O*).

DETAILED DESCRIPTION OF THE INVENTION CERAMIC SUPERCONDUCTORS

Superconducting oxides were first reported in 1964, but until recently, the intermetallic compounds showed higher superconducting temperatures. In 1975, research scientists at E. I. DuPont de Nemours discovered superconductivity in the system BaPb_(1-x) Bi_(x) O₃ with a Tc of 13 K. (A. W. Sleight et al., Solid State Commun., 17, 27 (1975)). The structure for the superconducting composition in this system is only slightly distorted from the ideal cubic perovskite structure. It is generally accepted that a disproportionation of the Bi(IV) occurs, namely, 2Bi(IV)(6s¹)→Bi(III)(6s²)+Bi(V)(6s⁰) at approximately 30 percent Bi. Sleight et al. found that the best superconductors were single phase prepared by quenching from a rather restricted single-phase region, and hence these phases are actually metastable materials. At equilibrium conditions, two phases with different values of x would exist; the phase with a lower value of x would be metallic and with a higher value of x would be a semiconductor. It is important to keep in mind that the actual assignment of formal valence states is a convenient way of electron accounting; the actual states include appreciable admixing of anion functions. Recently, for example, Cava and Batlogg, Nature, 332, 814 (1988), have shown that Ba₀.6 K₀.4 BiO₃ gave a Tc of almost 30 K., which is considerably higher than the 13 K. reported for BaPb₀.75 Bi₀.25 O₃.

La₂ CuO₄ was reported by Longo and Raccah, J. Solid State Chem., 6, 526 (1973), to shown an orthorhombic distortion of the K₂ NIF₄ structure with a=5.353Å, b=5.409Å and c=13.17Å. It was also reported that La₂ CuO₄ has a variable concentration of anion vacancies and may be represented as La₂ CuO_(4-x). Superconductivity has been reported for some preparations of La₂ CuO₄. See D. C. Johnston et al., Phys. Rev. B, 36, 4007 (1987). However, there appears to be some questions as to the stoichiometry of these products since only a small portion of the material seems to exhibit superconductivity (P. M. Grant et al., Phys. Rev. Letters, 58, 2482 (1987)).

The La_(2-x) A_(x) CuO₄ ceramics (A=CA, Sr, Ba), exhibit a tc of about 40 K. The substitution of the alkaline earth cation for the rare earth depresses the tetragonal-to-orthorhombic transition temperature. The transition disappears completely at x>0.2, which is about the composition for which superconductivity is no longer observed.

The compound YBa₂ Cu₃ O_(7-x) shows a superconducting transition of about 93 K. and crystallizes as a defect perovskite. The unit cell of YBa₂ Cu₃ O_(7-x) is orthorhombic (Pmmm) with a=3.8198(1)Å, b=3,8849(1)Å and c=11.6762(3)Å. The structure may be considered as an oxygen-deficient perovskite with tripled unit cells due to Ba-Y ordering along the c-axis. For YBa₂ Cu₃ O_(7-x), the oxygens occupy 7/9 of the anion sites. One-third of the copper is in four-fold coordination and two-thirds are five-fold coordinated. A reversible structural transformation occurs with changing oxygen stoichiometry going from orthorhombic at x=7.0 to tetragonal at x=6.0 (see P. K. Gallagher et al., Mat. Res. Bull., 22, 995 (1987)). The value x=7.0 is achieved by annealing in oxygen at 400°-500° C. and this composition shows the sharpest superconducting transition.

Recently, Maeda et al., cited above, reported that a superconductive transition temperature of 120 K. was obtained for Bi₂ CaSr₂ Cu₂ O₈. In most of the studies reported to date on the Bi/Ca/Sr/Cu/O system, measurements were made on single crystals selected from multiphase products. The group at DuPont selected platy crystals having a composition Bi₂ Sr_(3-x) Ca_(x) Cu₂ O_(8+y) (0.9>×>0.4) which showed a Tc of about 95 K. Crystals of Bi₂ Sr_(3-x) Ca_(x) Cu₂ O_(8+y) for x=0.5 gave orthorhombic cell constants a=5.399, b=5.414, c=30.904 (M. A. Subramanian et al., Science, 239, 1015 (1988)). The structure consists of pairs of CuO₂ sheets interleaved by Ca(Sr), alternating with double bismuth-oxide layers.

There are now three groups of superconducting oxides which contain the mixed Cu(II-Cu(III) oxidation states, namely La_(2-x) A_(x) CuO₄ where A=Ba, Sr or Ca; RBa₂ Cu₃ O_(7-x) where R is almost any lanthanide; and Bi₂ Sr_(2-x) Ca_(1-x) Cu₂ O_(8+y). Z. Z. Sheng and A. M. Herman, Nature, 332, 55 (1988) have recently reported on a high-temperature superconducting phase in the system Tl/Ba/Ca/Cu/O. Two phases were identified by R. M. Hazen et al., Phys. Rev. Lett., 60, 1657 (1988) namely Tl₂ CaBa₂ Cu₂ O₉ and Tl₂ Ba₂ Ca₂ Cu₃ O₁₀. A. W. Sleight et al. have also reported on the structure of Tl₂ Ba₂ CaCu₂ O₈ as well as Tl₂ Ba₂ CuO₆ [M. A. Subramanian and A. W. Sleight et al., Nature, 332, 420 (1988); J. B. Parise and A. W. Sleight et al., J. Solid State Chem., 76, 432 (1988)]. In addition, superconductor Tl₂ Ba₂ Ca₂ Cu₃ O₁₀ has been prepared by the DuPont group and shows the highest Tc of any known bulk superconductor, namely, 125 K.

A series of oxides with high Tc values has now been studies for the type (A^(III) O)₂ A₂ ^(II) Ca_(n-1) Cu_(n) O_(2+2n), where A(III) is Bi or Tl, A(II) is Ba or Sr, and n is the number of Cu-O sheets stacked. To date, n=3 is the maximum number of stacked Cu-O sheets examined consecutively. There appears to be a general trend whereby Tc increases as n increases.

Preparation

The present method can be used for the passivation of superconductive materials, regardless of the method of their preparation. Superconductive ceramic oxides are generally prepared by heating an intimate mixture of the oxide or carbonate powders of the solid elements at temperatures between 900° C. and 1100° C. After regrinding and reheating, the mixture is pressed into pellets and sintered (bonded without melting) at high temperatures for several hours. The pellets are then annealed at a lower temperature, optimally 400° C. to 450° C.

The conditions under which the ceramic is prepared affect its oxygen content. This is important because structural studies have shown that the number and arrangement of oxygen atoms in the lattice is important to the oxide's superconductive properties. In fact, to produce the highest superconducting transition temperatures, the ceramic is preferably heated in an atmosphere of pure oxygen. The ceramic should then be cooled slowly for about 5-6 hours in a furnace.

Characterization of Superconductor Surfaces and the Overlayer-Substrate Interface

Electron spectroscopy was used to identify occupied and unoccupied electronic states of the superconductors La₁.85 Sr₀.15 CuO₄, YBa₂ Cu₃ O_(7-x), Bi₂ Ca_(1+x) Sr_(2-x) Cu₂ O_(8+y), and related cuprate compounds. Results from polycrystalline and single crystalline materials reveal that the valence band emission for the 2-1-4, 1-2-3, 2-1-2-2 and 2-2-2-3 superconductors are remarkably similar, with a low density of states near E_(f), a central feature at about 3.4 eV, and shoulders at 2.1 and 5.4 eV that are derived primarily from Cu-O hybrid states. Comparison with calculated densities of the states shows the importance of correlation effects. Structure in the empty states can be related to La 5d, La 4f, Ba 5d, Ba 4f, Y 4d, Ca 3d, Sr 4d, Bi5p and O 2p empty states. Core level results show the Cu 2p_(3/2) main line and satellite structure associated with formal Cu²⁺ configuration. The O is emission reveals inequivalent chemical environments. Interface studies can show substrate disruption, contact formation, and passivation, depending on the overlayer. Ag and Au overlayers have a minimal effect (inert contacts) while the deposition of oxygen-scavenging atoms (Ti, Fe, Cu, Pd, La, Al, In, Bi and Ge) results in oxygen removal and surface disruption.

The invention will be further described by reference to the following detailed example wherein x-ray photoelectron spectroscopy (XPS) measurements of the interface formation were performed under standard ultrahigh vacuum conditions (pressure during measurement less than 1×10⁻¹⁰ Torr). See S. A. Chambers et al., Phys. Rev. B, 35, 634 (1987). a monochromatic beam of Al k.sub.α photons (hv=1486.6 eV) was focused on the sample surface and the energy of the emitted electrons was measured with a Surface Science Instruments hemispherical analyzer using a resistive anode position-sensitive detector. The photoelectrons were collected at an angel of 60° relative to the surface normal and the cone of acceptance of the analyzer was 30°. The x-ray beam diameter was 300 μm at the sample and the pass energy of the analyzer was 50 eV. The Cu 2p_(3/2) emission from pure Cu was used to calibrate the binding energy scale (binding energy of 932.5 eV). Data were acquired and analyzed on a dedicated HP9836C computer. For each of the interfaces studies, the XPS energy distribution curves (EDCs) were measured for the valence bands and the Cu 2p and O ls core levels, as well as Y 3d, Ba 3d, Bi 4f, Ca 2p, and Sr 3d core levels, as appropriate. Core level binding energy shifts and the attenuation of the substrate core level emission as a function of overlayer thickness were used to identify reacting species and overlayer growth morphology.

YBa₂ Cu₃ O_(7-x) was obtained from Argonne National Laboratory, Ill. Clean surfaces of the various substrate materials for these tests were prepared by fracturing the samples in situ, which exposed uncontaminated internal surfaces. Prior to deposition of any material, the clean surfaces were thoroughly characterized to allow comparison of the initial surfaces and those which had been coated with a particular overlayer.

Deposition of oxides on the superconductors was performed by evaporating the parent material (e.g., Bi, Al or Si) in an ambient atmosphere of activated oxygen. Bi and Si were evaporated from resistively-heated Ta boats, and Al was evaporated from a resistively-heated W basket. Before any deposition, the evaporators were thoroughly degassed (pressure in the chamber was less than 2×10⁻¹⁰ Torr during evaporation, prior to introducing O₂), and the evaporation rates were stabilized at a rate of about 1 Å/minute at the sample. The source-to-sample distance was about 30 cm, and the evaporation rate was measured with an Inficon quartz crystal thickness monitor. High-purity O₂ from a stainless steel gas mainfold could be introduced to the chamber via a leak-valve, and the pressure monitored with an ionization gauge. Composition of the gases in the chamber was determined with a quadupole mass spectrometer. The partial pressure of O₂ in the deposition chamber was 1.0×10⁻⁶ Torr, and the partial pressure of the residual gases (predominantly CO and H₂ O) was less than 1×10⁻⁹ Torr. Between the evaporator and the sample was a thorium oxide-coated iridium filament which was electrically isolated from the chamber walls. Passing current through this filament and electrically biasing it negatively relative to the chamber provided a source of energetic electrons to dissociate O₂ molecules and ionize O atoms. This activation process had the effect of increasing their chemical reactivity.

Passivation layers were obtained using the following process. First, the sample was cleaved and positioned facing away from the evaporator. Second, a previously degassed boat containing the desired parent metal or CaF₂ was heated to the evaporation temperature and the rate of evaporation stabilized. Oxygen was then leaked into the chamber to the desired pressure. (The addition of O₂ did not affect the evaporation rates, and other experiments have shown that evaporation of these materials in a non-activated O₂ ambient did not produce any significant oxide formation on the substrate, except that produced by reaction with oxygen atoms from the substrate. It is these reactions with the substrate which the activated deposition process inhibits.)

Once the O₂ pressure in the chamber was stabilized, the filament was heated and biased to -250 V to establish a 50 mA emission current between filament and ground. Mass spectroscopy of the gases present at this point showed O atoms and O₂ molecules in equal proportions (partial pressure=5×10⁻⁷ Torr) and doubly ionized O (O⁻²) present at roughly 10% of the total pressure (partial pressure=1×10⁻⁷ Torr). The sample was then turned to face the evaporator to deposit a particular thickness of oxide. After each deposition, XPS spectra were taken to determine changes on the surface. The Cu 2p_(3/2) lineshape was used as an indicator of reaction and to evaluate the emission ratio of the satellite to the main line. The O ls lineshape was used to qualitatively confirm deposition of oxide layers.

Vapor deposition of CaF₂ was done in the same system, without the addition of oxygen using vacuum evaporation technique. Solid CaF₂ was evaporated from tungsten boats, and was deposited on the substrate surface.

EXAMPLE Deposition of Dielectric Layers on YBa₂ Cu₃ O_(7-x) A. Deposition of Bi-O and CaF₂

A direct measure of the chemical reactivity and modification of the superconductor surface region can be gained by following the behavior of the Cu 2p_(3/2) core level emission feature using x-ray photoemission. The bottom-most curve of FIG. 1 shows the Cu 2p_(3/2) lineshape for the clean surface of a freshly fractured sample of YBa₂ Cu₃ O_(7-x). This complex lineshape, which is observed for the 1-2-3 materials as well as the 2-1-4 and 2-1-2-2 superconductors, indicates that the Cu valence is nominally 2+. We have found that any modification of the superconductor by reactive metal deposition or Ar ion bombardment leads to the loss of the doublet features at higher binding energy, termed the "satellite." The loss of this feature indicates a valence change to Cu¹⁺, structural modification of the lattice, and the loss of superconductivity.

The top-most spectrum of FIG. 1 shows that the deposition of only 2.5 Å of a reactive metal, Cu, leads to the almost complete loss of satellite doublet (satellite-to-mainline-intensity ratio reduced from 0.34 for the clean surface to 0.04). Such effects are typical for reactive metals, and to date, Ti, Fe, Cu, Pd, La, Al and In have been examined, together with the semiconductors Ge and Si. The results indicate that the Cu valence has changed within a distance of at least 50-60 Å of the surface and that a nonsuperconducting, insulating layer is formed, i.e., a poor electrical contact. Moreover, the amount of deterioration will be enhanced by thermal processing. The only metal which leads to no change in the satellite-to-mainline-intensity ratio is Au. The deposition of adatoms of Au leads to a metallic overlayer that does not adversely effect the substrate. The deposition of 10 Å of CaF₂ also leaves the substrate completely intact with no evidence of disruptive interactions. Moreover, CaF₂ appears to cover the surface uniformly, as judged by the exponential attenuation of the substrate emission. In contrast to the other overlayers discussed hereinabove, CaF₂ is a large bandgap insulator with a high static dielectric constant and it is useful as an insulating layer in device fabrication.

As part of these passivation studies, the deposition of metals in an activated-oxygen environment at 1×10⁻⁶ Torr partial pressures of oxygen was examined. In this way, it was possible to provide activated oxygen from the gas phase to form metal oxide precursors that would not react with oxygen from the superconductor.

In FIG. 1, the results for Bi deposited onto Yba₂ Cu₃ O_(7-x) both with an without activated oxygen are compared. While not as reactive as Cu metal, an 8 Å layer of Bi leads to disruption/reaction on Yba₂ Cu₃ O_(7-x) as discussed above. In contrast, the Cu 2p lineshape did not change following the deposition of 6 Å of Bi in activated oxygen, indicating that the Bi-O overlayer did not significantly modify the superconductor surface.

B. Deposition of SiO₂ and Al-W Oxides

FIGS. 2 and 3 summarize the results obtained for the deposition of SiO₂ and Al-W oxide layers by the activated chemical vapor deposition of Si and Al. In FIG. 3, the O ls spectrum for the uncoated ("clean") substrate exhibits a wide asymmetric main peak, indicating the presence of more than one chemical environment. It is made up of a dominant component at about 529 eV, derived from Cu-O planes, and a shallower component, attributed to Cu-O chains. The deposition of Al from a tungsten basket in an activated oxygen atmosphere (filament bias was 250 eV; filament emission was 50 mA; O₂ pressure=1×10-5 Torr) resulted in the deposition of an overlayer that is believed to comprise an about 1:1 ratio of aluminum:tungsten oxides. The non-disruptive nature of these dielectric oxide overlayers is confirmed by the persistent emission from the Cu 2p_(3/2) satellite structure at about 942 eV binding energy in FIG. 2. FIG. 3 shows the corresponding O_(ls) core level emission, which demonstrates that new oxygen bonding configurations are observed following deposition.

C. Deposition of Al₂ O₃

Evaporation was done from a coiled W basket in an ambient pressure of 1×10⁻⁶ Torr O₂. The emission current from the thoria-coated filament was 30 mA, and the biasing voltage was -250 V. Apparently, the lower pressure and activation current in this test did not create detectable amounts of the volatile W oxide on the basket.

FIG. 4 demonstrates the Cu 2p_(3/2) peak for the deposition of 16 Å pure Al (dot-dash-line) and for 16 Å Al+activated O₂ (solid line). The integrated intensity of the satellite emission for the pure Al case was 12% of the main peak intensity, compared to 22% for the activated deposition. The main peak for the activated deposition is also broader than for the pure Al deposition, although the total integrated intensity was the same. Both of these results imply that the deposition of pure Al converts more Cu within the probe depth from the 2+ oxidation state of the superconductor to the 1+ state, characteristic of the disrupted superconductor, and removes more oxygen from the superconductor. These results verify that the deposition of aluminum oxide reacted less with the substrate and withdrew less oxygen from the superconductor than pure Al. No W was detected on the surface after this test.

The invention has been described with reference to various specific and preferred embodiments and techniques. However, it should be understood that many variations and modifications can be made while remaining within the spirit and scope of the invention. 

What is claimed is:
 1. A method comprising evaporating Al, Bi, Si or Al-W in an atmosphere of activated oxygen in the presence of a superconductive ceramic oxide surface so that a layer of an oxide of Al, Bi, Si or Al-W is applied to said surface, wherein said oxide layer is effective to passivate said superconductive oxide surface without disrupting the superconductive properties.
 2. The method of claim 1 wherein said oxide passivation layer is applied to the surface of a 2-1-4 superconductive ceramic oxide.
 3. The method of claim 1 wherein said oxide passivation layer is applied to the surface of a 1-2-3 superconductive ceramic oxide comprising Y.
 4. The method of claim 1 wherein said oxide passivation layer is applied to the surface of a 2-1-2-2 superconductive ceramic oxide comprising Bi.
 5. The method of claim 1 wherein said oxide passivation layer is applied to the surface of a 2-1-2-2 superconductive ceramic oxide comprising Tl.
 6. The method of claim 1 wherein said oxide passivation layer is applied to the surface of a superconductive ceramic oxide of the formula RBa₂ Cu₃ O_(7-x), wherein R is a lanthenide element.
 7. The method of claim 1 wherein the layer is at least about 50 Å in thickness. 